Structural and surface characterizations of 2D β-In2Se3/3D β-Ga2O3 heterostructures grown on c-Sapphire substrates by molecular beam epitaxy

Integrating two-dimensional (2D) layered materials with wide bandgap β-Ga2O3 has unveiled impressive opportunities for exploring novel physics and device concepts. This study presents the epitaxial growth of 2D β-In2Se3/3D β-Ga2O3 heterostructures on c-Sapphire substrates by plasma-assisted molecular beam epitaxy. Firstly, we employed a temperature-dependent two-step growth process to deposit Ga2O3 and obtained a phase-pure \documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$(\overline{2 }01)$$\end{document}(2¯01) β-Ga2O3 film on c-Sapphire. Interestingly, the in-situ reflective high-energy electron diffraction (RHEED) patterns observed from this heterostructure revealed the in-plane ‘b’ lattice constant of β-Ga2O3 ~ 3.038Å. In the next stage, for the first time, 2D In2Se3 layers were epitaxially realized on 3D β-Ga2O3 under varying substrate temperatures (Tsub) and Se/In flux ratios (RVI/III). The deposited layers exhibited (00l) oriented β-In2Se3 on \documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$(\overline{2 }01)$$\end{document}(2¯01) β-Ga2O3/c-Sapphire with the epitaxial relationship of \documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$[11\overline{2 }0]$$\end{document}[112¯0] β-In2Se3 || [010] β-Ga2O3 and \documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$[10\overline{1 }0]$$\end{document}[101¯0] β-In2Se3 || [102] β-Ga2O3 as observed from the RHEED patterns. Also, the in-plane ‘a’ lattice constant of β-In2Se3 was determined to be ~ 4.027Å. The single-phase β-In2Se3 layers with improved structural and surface quality were achieved at a Tsub ~ 280 °C and RVI/III ~ 18. The microstructural and detailed elemental analysis further confirmed the epitaxy of 2D layered β-In2Se3 on 3D β-Ga2O3, a consequence of the quasi-van der Waals epitaxy. Furthermore, the β-Ga2O3 with an optical bandgap (Eg) of ~ 5.04 eV (deep ultraviolet) when integrated with 2D β-In2Se3, Eg ~ 1.43eV (near infra-red) can reveal potential applications in the optoelectronic field.


RHEED
Since the advent of two-dimensional (2D) layered van der Waals (vdWs) materials, Indium Selenide (In x Se y ), one of the prominent candidates in this family, has been widely explored in the scientific community to investigate its novel and impeccable properties [1][2][3][4][5][6][7][8] .It belongs to a complex system that crystallizes into different stoichiometric ratios (stacking configurations), e.g., InSe (β, γ), In 2 Se 3 (α, β, γ and δ), In 3 Se 4, etc., under various deposition conditions and growth techniques 4,7,8 .Among these, the β-In 2 Se 3 with its rhombohedral crystal structure has a primitive unit cell (a = b = 4.00 Å, and c = 28.33Å) comprising three monolayers that are stacked vertically and repeatedly by weak vdW forces, with an in-plane covalently bonded atomic sequence of "Se − In − Se − In − Se" 2 .The β-In 2 Se 3 is renowned for its exceptional chemical stability and remarkable optical activity at room temperature (RT) and further exhibits strong 2D quantum confinement effects with its absorption edge in the near infra-red (IR) spectral range (~ 1.43 eV) 1,2 .In addition, recent studies demonstrated that the phase-engineering of In 2 Se 3 from α to β through thermal annealing has resulted in ultrahigh responsivity and detectivity of 8.8 × 10 4 A/W and 2.9 × 10 13 Jones, respectively 6 .
On the other hand, recently, there has been significant research interest in exploring the integration of 2D layered materials with wide bandgap (WB) semiconductors, particularly Gallium Oxide (Ga 2 O 3 ) [9][10][11][12][13] .Being a fourth-generation semiconducting material, Ga 2 O 3 , one of the group-III metal sesquioxide exhibits various polymorphs: α, β, γ, δ, and ε 14 .Among which the monoclinic β-phase (a = 12.23 Å, b = 3.04 Å, c = 5.80 Å, and β = 103.71°)with its direct bandgap (E g ) ~ 4.9 eV is considered to be the thermodynamically stable structure 14,15 .Due to its ultrawide E g , high breakdown electric field of ~ 8 MV/cm, and robust chemical/thermal stability, it has exhibited tremendous progress in high-power electronics and deep ultraviolet (UV) optoelectronic devices [16][17][18] .Integrating this material with 2D layered materials can unveil novel opportunities in device physics.For instance, Wang et al. demonstrated a solar-blind photodetector with p-GaSe/n-Ga 2 O 3 vdWs heterostructure that showed a high responsivity of 51.9 A/W and a pronounced specific detectivity up to 10 14 Jones, resulting from the efficient separation of charge carriers across the pn junctions 11 .An ambipolar p-TMD (p-MoTe 2 or p-WSe 2 )/n-Ga 2 O 3 junction field effect transistor (JFET) was reported by Choi et al., with two different types of channels in a single device architecture with their respective charge carriers 12 .Despite the challenge in realizing the enhancement mode (e-mode) operation of the Ga 2 O 3 device due to lack of p-type doping, Yang et al. fabricated a β-Ga 2 O 3 FET with ferroelectric α-In 2 Se 3 wrapped-gate that changed from depletion-to e-mode operation by effectively controlling the threshold voltage 13 .These findings collectively highlight the significance and potential of 2D material/β-Ga 2 O 3 heterostructures for future device applications.
Nevertheless, the integration of these 2D materials with WB-Ga 2 O 3 from the previous works was constrained to the ex-situ techniques, particularly by exfoliation or transfer methods of either the 2D vdW layers or the underlying Ga 2 O 3 layers from the bulk substrates.Although the results are encouraging, such methods offer limited control over film thickness, may be prone to contamination and defects, and, most importantly, accessible with reduced scalability, therefore limiting their usage in large-area applications.Owing to these challenges, the utilization of molecular beam epitaxy (MBE) emerges as a proven growth technique to fabricate these heterostructures in situ with its ultra-high vacuum (UHV) conditions, high pure elements, thickness controllability and further yielding single crystalline materials with reduced defects.
For the first time in this study, the mixed-dimensional 2D β-In 2 Se 3 /3D β-Ga 2 O 3 heterostructures were realized in situ using plasma-assisted molecular beam epitaxy (PA-MBE) on c-Sapphire.To achieve high-quality heteroepitaxial films, careful optimization of the initial β-Ga 2 O 3 growth process is essential.A strategic approach involves the introduction of a low-temperature (LT) buffer (nucleation) layer, which proves effective in two key aspects.Firstly, the LT buffer layer serves as a sacrificial template by incorporating and localizing threading dislocations (TDs) that arise due to the lattice mismatch concerning the substrate 19,20 .Secondly, it provides a homo-surface, circumventing lattice constraints 21 and facilitating a smoother transition for high-temperature (HT) film growth.Consequently, we used a two-stepped β-Ga 2 O 3 film grown under LT and HT conditions, commonly used for the heteroepitaxy on a Sapphire substrate, to improve the crystal quality effectively 19,21 .Amidst the daunting challenge of the inherent and uneven surface of 3D Ga 2 O 3 , we successfully achieved the epitaxy of 2D In 2 Se 3 , thanks to our vigilant in-situ reflective high-energy electron diffraction (RHEED) tool for providing the information about the structural changes, in-plane lattice constants, and epitaxial relationships of the grown films.Besides the rich phases of In 2 Se 3, we achieved the dominant phase 2D β-In 2 Se 3 on 3D β-Ga 2 O 3, which was confirmed by X-ray Diffraction (XRD) and Raman Spectroscopy.Furthermore, the surface morphological changes of the grown layers were studied carefully using Atomic Force Microscope (AFM) measurements.The microstructural and detailed elemental analysis across the heterostructures grown on c-Sapphire was thoroughly investigated by (Scanning) Transmission Electron Microscopy-(S)TEM measurements.The results presented in this study establish a fundamental understanding of the epitaxy of 2D In 2 Se 3 /3D Ga 2 O 3 heterostructures, which is crucial for its commercialization in large-area applications.

Results and discussion
Figure 1a,b shows the in-situ RHEED patterns of the c-Sapphire substrate (before growth).Soon after the growth of LT-Ga 2 O 3 film (substrate temperature, T sub ~ 450 °C), the transition in the RHEED patterns occurred along both the azimuthal directions (repeated for every 60° rotation), as shown in Fig. 1c,d  This preferential alignment of β-Ga 2 O 3 is attributed to its similar oxygen atomic arrangement compared to c-Sapphire, and the effort to minimize lattice strain to establish a coherent epitaxial relationship between them.This is attributed to the fact that the lattice points within the growth directional planes, (0001) Sapphire and (201) β-Ga 2 O 3, maintain closer repeated interatomic distances (rectangle with white circled corners), as shown in Fig. 1g,h 038 Å, which matches well with the theoretical value 15,24 , which could indicate the fully relaxed β-Ga 2 O 3 film grown on c-Sapphire.This was determined quantitatively by the translation of streak spacing in the reciprocal lattice by the number of pixels achieved from the RHEED patterns 26 .The additional diffraction streaks observed in this direction (indicated by blue arrows) also maintained a similar streak spacing.This coexistence of patterns along [102] β-Ga 2 O 3 might arise from the octahedral and tetrahedral planes of Ga atoms within the (201) β-Ga 2 O 3 ; however, further understanding is required to confirm its origin.The surface morphology of the as-grown LT-Ga 2 O 3 and two-stepped Ga 2 O 3 films grown on c-Sapphire is shown in Fig. 2b,c from the AFM scans.At LT, the surface of the film exhibited small granular morphology with dense grain boundaries owing to a root mean square (RMS) of ~ 0.56 nm.This is because the adatoms at LT will not have sufficient energy to transfer and nucleate with adjacent atoms on the surface, thus resulting in high nucleation sites, as seen in Fig. 2b.On the other hand, this decreased mobility of surface species can promote uniform dispersion of nuclei that can effectively cover the substrate 19 and provide a homo-surface for the HT film growth.At HT conditions, Fig. 2c, the surface of the film was covered with large grains with reduced grain boundaries and exhibited a rougher surface (RMS ~ 5.83 nm).This could be ascribed to the greater likelihood of an adatom encountering an existing island formed during the ripening stage and promoting further growth primarily due to an increased adatom diffusion coefficient at HT.
Figure 3a shows XRD 2θ-scans of the Ga 2 O 3 films grown on c-Sapphire.The LT nucleation film exhibited distinct diffraction peaks at ~ 18.9° and ~ 38.3°, corresponding to the (201) and (402) diffraction peaks of β-Ga 2 O 3 , respectively 15,17 (ICDD Card No. 01-082-3838).Moreover, the ratio calculated between these (201) and (402) diffraction peaks was found to be ~ 1.89 (ideal value ~ 2.2) 27 , suggesting the dominant β-Ga 2 O 3 when grown at a low T sub ~ 450 °C, in contrast to the observation of secondary phases by Oshima et al. 27 Also, the d-spacing measured between the (201) planes of LT β-Ga 2 O 3 is determined to be ~ 0.468 nm, with the thickness of film ~ 32 nm as shown by TEM images in the supplementary information, Fig. S3.Further, depositing the HT Ga 2 O 3 layers, ) relates to the stretching and bending of GaO 4 tetrahedra [28][29][30] .The more pronounced vibrational modes were observed from the two-stepped film with the FWHM of A g 3 mode ~ 5.5 cm −1 .Figure 3c shows the transmittance spectra of the HT/LT β-Ga 2 O 3 film at RT, and the spectrum exhibited an average transmittance of ~ 94% with clear interference fringes in the visible region.The direct bandgap of a semiconductor can be determined from the UV-visible spectra by using the Tauc relation (αhν) 2 ∝ (hν − E g ) 31 , where α is the absorption coefficient, hν is the incident photon energy, and E g is the optical bandgap.An abrupt decrement in the wavelength was observed at the absorption edge around 250 nm, indicating the presence of an optical bandgap.The E g of the β-Ga 2 O 3 film is estimated by extrapolating the intercept of the energy axis at α = 0 with an approximate value ~ 5.04 eV (deep UV region), similar to previously reported values 30,32,33 .By virtue of the successful epitaxy of two-stepped β-Ga 2 O 3 film on c-Sapphire, in the following crucial stage, we focused on depositing and investigating the 2D-In 2 Se 3 films on 3D β-Ga 2 O 3 using MBE.
The epitaxy of the chalcogenide material in a typical solid source UHV-MBE system is employed at a high chalcogen-to-metal flux ratio due to increased volatility and lower sticking coefficient of chalcogen atoms at the growth surface 34 .Owing to this, we maintained Se-rich conditions for the In 2 Se 3 epitaxy in this work by setting the Se/In flux ratio (R VI/III ) larger than ~ 15 8 .Figure 4a,b shows the RHEED patterns of the as-grown HT/LT β-Ga 2 O 3 /c-Sapphire before the growth of In 2 Se 3 .When the growth was maintained at T sub of 480 °C (R VI/III ~ 28), the RHEED patterns remained similar to β-Ga 2 O 3 for the whole growth, as seen in Fig. 4c,d.This indicates that the epitaxy of In 2 Se 3 layers didn't occur at these conditions, which might be caused by the kinetic limitations (rate of adsorption and desorption of adatoms) or nucleation barriers 35 that may not be favorable at 480 °C, leading to hindered growth.Further decreasing the T sub ~ 330 °C, the transition in the RHEED patterns was observed along both the azimuthal directions, as shown in Fig. 4e,f, indicating the change in the crystal structure from monoclinic β-Ga 2 O 3 to rhombohedral β-In 2 Se 3 structure.In addition, the streak spacing ratio between the a-a/m-m planes was measured to be ~ √ 3 , representing the six-fold symmetry of In 2 Se 3 layers.Despite the 3D surface morphology of the β-Ga 2 O 3 film, the In 2 Se 3 layers were successfully grown on it.This could be a consequence of quasi-vdWs epitaxy being independent of the surface lattice conditions of the underlying layer 36 .The In 2 Se 3 layers grown on β-Ga 2 O 3 film followed the in-plane epitaxial relationship of  On the other hand, it has been a challenging issue to identify and differentiate the commonly obtained crystal phases of In 2 Se 3 , among which the rhombohedral crystal structures of α-and β-In 2 Se 3 share similar but different space groups (R3m and R 3m) 5,38 .The primary difference between these two structures lies in the location of In atoms at tetrahedral and octahedral sites covered by the Se packing in α-and β-In 2 Se 3, respectively 38 .Subsequently, insisting on a demanding characterization method and prudent analysis to distinguish the respective crystal phases.Figure 6a 6c, which are characteristics of β-In 2 Se 3 , that are similar to the previously reported results 1,2,39 .A similar peak overlapping was observed between the pronounced A g 3 mode from the β-Ga 2 O 3 and the A 1 (LO) mode of β-In 2 Se 3, as shown in Fig. S6(c,d).The active vibrational modes exhibited by β-In 2 Se 3 with regard to the similarly structured α-In 2 Se 3 are validated upon comparing the typical Raman peaks of α-In 2 Se 3 as summarized by Liu et al. 39 .However, the sample grown at T sub of 280 °C under the R VI/III ~ 38 exhibited two additional Raman modes ~ 151.3 cm −1 and ~ 252.6 cm −1 , along with the respective β-In 2 Se 3 modes.These peaks are characteristics of γ-In 2 Se 3 , with the former phonon mode corresponding to the zone center vibration and the latter to the excess contribution of Se atoms' linkage to the Se-Se bond due to the high Se flux used in this series 40,41 .This indicates that, at these epitaxial conditions, the growth leads to the co-existence of β-In 2 Se 3 and γ-In 2 Se 3 with the dominance in the former phase.The existence of additional 3D γ-In 2 Se 3 may cause predominantly vertical growth, resulting in a rougher surface, as evident from the surface morphology characterization mentioned above.
Furthermore, Fig. 7a shows the STEM high-angle annular dark-field (HAADF) cross-sectional view of the β-In 2 Se 3 /β-Ga 2 O 3 heterostructure grown on c-Sapphire.The thicknesses of the LT-, HT β-Ga 2 O 3, and β-In 2 Se 3 films are determined to be ~ 32 nm, ~ 120 nm, and ~ 28 nm, respectively, with the corresponding growth rates of ~ 0.53, ~ 1.0, and ~ 0.47 nm/min.Figure 7b,c provides a detailed visualization of the interfaces between the LT β-Ga 2 O 3 /c-Sapphire and β-In 2 Se 3 /HT β-Ga 2 O 3 heterostructures.Regardless of the relatively rough surface of the two-stepped β-Ga 2 O 3 film that may result in a non-abrupt 2D/3D interface, the layered structure of β-In 2 Se 3 is clearly observed, attributed to the quasi-van der Waals epitaxy.The detailed elemental mappings of the entire β-In 2 Se 3 /β-Ga 2 O 3 heterostructure grown on c-Sapphire are shown in Fig. 7d-i, which reveal an abrupt transition and uniform distribution of respective elements within the specific layers.These results evidently support the objective of the present work on realizing the epitaxial growth of β-In 2 Se 3 on β-Ga 2 O 3 , which shows the potential scope to study the mixed dimensional heterostructures for future device applications using MBE technique.
So far, we have successfully achieved the epitaxy of the single phase 2D β-In 2 Se 3 layers on 3D two-stepped β-Ga 2 O 3 /c-Sapphire.In the present series, the better structural and surface quality of β-In 2 Se 3 layers was obtained when grown at R VI/III and T sub of ~ 18 and ~ 280 °C, respectively.Finally, such a mixed dimensional (2D β-In 2 Se 3 /3D β-Ga 2 O 3 ) heterostructure can avail the benefits offered by both materials, specifically in the optoelectronic field, with its absorption edges extending from Near-IR (~ 1.43 eV) 1 to deep UV regions (~ 5.04 eV), and can be used as a dual-band photodetector.Also, the epi-grown In 2 Se 3 and Ga 2 O 3 , being intrinsically n-type, can form a heterostructure exhibiting an nN isotype heterojunction, forming 2DEG upon bandgap engineering.
Figure 1a,b shows the in-situ RHEED patterns of the c-Sapphire substrate (before growth).Soon after the growth of LT-Ga 2 O 3 film (substrate temperature, T sub ~ 450 °C), the transition in the RHEED patterns occurred along both the azimuthal directions (repeated for every 60° rotation), as shown in Fig. 1c,d, indicating the change in crystal structure from rhombohedral (α) c-Sapphire to monoclinic (β) Ga 2 O 3 .The in-plane epitaxial relationship observed from RHEED patterns revealed that the Ga 2 O 3 was aligned along [010] β-Ga 2 O 3 || [1010] c-Sapphire and [102] β-Ga 2 O 3 || [1120] c-Sapphire, and the respective growth directional views were showed in Fig. S1.

Figure 1 .
Figure 1.In-situ RHEED patterns of (a,b) c-Sapphire substrate (before growth), (c,d) LT β-Ga 2 O 3, and (e,f) HT/LT β-Ga 2 O 3 films at ~ 5 min of their respective growths.The in-plane crystallographic views of (g) (0001) Sapphire and (h) (201) β-Ga 2 O 3 along the growth direction were visualized using the ball and stick model by VESTA Software 42 .
. Further exerting an in-plane lattice mismatch of -3.2% and -10.7% (minus indicates compressively strained β-Ga 2 O 3 ) in their respective directions.This mismatch might arise due to the slight distortion of regular hexagon redistribution of oxygen atoms with a single O-O distance: 4.76 Å in (0001) Sapphire to two O-O distances: 4.96 Å and 5.15 Å in (201) β-Ga 2 O 315 , owing to the change in ionic radii of the Al 3+ (0.54 Å) to Ga 3+ (0.62 Å)22 .Moreover, this initially grown LT-Ga 2 O 3 can serve as a nucleation film between the c-Sapphire and the HT-Ga 2 O 3 (T sub ~ 700 °C) film by minimizing this lattice and in-plane thermal expansion (α) mismatches between (0001)-Sapphire (α s ~ 5 × 10 -6 K −1 )23 and (201) β-Ga 2 O 3 (α g ~ 7.8 × 10 -6 K −1 ) 24 .Corroborating this, the evolution of RHEED patterns from the LT β-Ga 2 O 3 film grown on c-Sapphire exhibited a significant improvement in the crystalline quality from starting to finishing growth at T sub ~ 450 °C, as shown in Fig S2.Furthermore, Fig. 1e,f shows sharper and streakier RHEED patterns observed from the HT-grown Ga 2 O 3 film.This could result from the lattice mismatch compensation and uniform nucleation in the LT-Ga 2 O 3 film, providing a decent surface for the homoepitaxy at HT.More detailed information on the temporal evolution of RHEED patterns from the two-stepped (HT/LT) β-Ga 2 O 3 film indicating the improvement in crystalline quality is discussed in the supplementary information.Notably, the information about the 'b' lattice constant of the β-Ga 2 O 3 crystal structure can be obtained from the RHEED pattern along [102] β-Ga 2 O 3 .In this regard, we extracted the RHEED intensity profiles along respective directions, as shown in Fig. 2a.The reciprocal lattice spacing of b * G [102] = 2.068 Å −1 extracted from known lattice distance of c-Sapphire along b S [1120] = 8.241 Å ( √ 3a S ) 25 in real space, has yielded the 'b
[1120] β-In 2 Se 3 || [010] β-Ga 2 O 3 and [1010] β-In 2 Se 3 || [102] β-Ga 2 O 3 .The spotty pattern observed in this condition may have originated from the 3D growth of In 2 Se 3 layers.Gradually, the RHEED patterns became streakier upon further reducing the T sub ~ 280 °C, indicating improved lateral growth, as shown in Fig. 4g,h.Following this, the In 2 Se 3 layers were grown with varied R VI/III of 38 and 18 at a T sub of 280 °C.The RHEED patterns became broader with slight spots for the sample grown at an increased flux ratio of 38, Fig. 4i,j, suggesting the declined surface quality.Among the whole series, a clear and streakier pattern was observed for the entire epitaxy when grown at R VI/III ~ 18 (T sub ~ 280 °C), suggesting the improved surface of the In 2 Se 3 layers, as shown in Fig. 4k,l.Furthermore, the inplane reciprocal streak spacing along [1010] β-In 2 Se 3 with the respective [102] β-Ga 2 O 3 has yielded the real space 'a' lattice constant of β-In 2 Se 3 , b I [1010] = 2π/b * I [1010] Å ~ 4.027 Å as shown Fig. 4m.The surface morphologies of In 2 Se 3 layers grown on 3D β-Ga 2 O 3 films at varying epitaxial conditions are shown in Fig. 5 from the 5 × 5 µm 2 AFM scans.At a T sub of 330 °C (R VI/III ~ 28), we can observe a high density of smaller triangles (~ 250 nm) with a pronounced vertical stacking, resulting in a 3D surface morphology of In 2 Se 3 (RMS ~ 13.70 nm) which is in correspondence with the observation of spotty RHEED patterns.Further reducing the T sub ~ 280 °C, the density of
shows the XRD 2θ-scans of In 2 Se 3 layers grown on two-stepped β-Ga 2 O 3 /c-Sapphire at different epitaxial conditions.The sample grown at T sub ~ 480 °C (R VI/III ~ 28) exhibits only the diffraction peaks related to β-Ga 2 O 3 , indicating the absence of In 2 Se 3 epitaxy which agrees well with the observed RHEED patterns.On the other hand, at all the other epitaxial conditions, apart from the diffraction peaks of β-Ga 2 O 3 , we can observe two distinct additional peaks, ~ 9.4° and ~ 28.5° corresponding to (003) and (009) diffraction planes of rhombohedral β-In 2 Se 3 crystal structure (ICDD Card No. 35-1056, space group R 3m).Furthermore,